High-strength aluminum alloy coatings, deformation layers and methods of making the same

ABSTRACT

A high-strength aluminum alloy coating. The coating includes aluminum, 9R phase, fine grains, nanotwins, stacking faults, and a solute capable of stabilizing the 9R phase, the fine grains, and the stacking faults. A method of making a high-strength aluminum alloy coating on a substrate. The method includes, depositing the constituents of an aluminum alloy on a substrate such that the deposit forms a high-strength aluminum alloy coating containing 9R phase, fine grains, nanotwins, and stacking faults. A high-strength deformation layer in and on a casting of an aluminum alloy containing 9R phase, fine grains, nanotwins, stacking faults, and a solute capable of stabilizing the PR phase, the fine grains, and the stacking faults. A method of making a high-strength deformation layer in and on a casting of an aluminum alloy by deforming the alloy such that deformation layer contains 9R phase, fine grains, nanotwins, and stacking faults.

CROSS-REFERENCE TO RELATED APPLICATIONS

The present U.S. patent application is a divisional of U.S. patentapplication Ser. No. 15/949,436 filed Apr. 10, 2018, which is related toand claims the priority benefit of U.S. Provisional Patent ApplicationSer. No. 62/484,771 filed Apr. 12, 2017, the contents of which arehereby incorporated by reference in their entirety into the presentdisclosure.

STATEMENT REGARDING GOVERNMENT FUNDING

This invention was made with government support under DE-SC0016337awarded by the Department of Energy. The government has certain rightsin the invention.

TECHNICAL FIELD

This disclosure relates high-strength aluminum alloy coatings onsubstrates and high-strength deformation layers in and on aluminum alloycastings with engineered microstructural and grain boundary features andmethods of making such high-strength aluminum alloy.

BACKGROUND

This section introduces aspects that may help facilitate a betterunderstanding of the disclosure. Accordingly, these statements are to beread in this light and are not to be understood as admissions about whatis or is not prior art.

The ever-increasing demand for high-strength, lightweightautomotive-grade materials, driven by the requirement for high fuelefficiency, makes development of novel ultra-strong and ductile Aluminum(Al) alloys a priority. Steels that hold analogous specific strengthoutperform Al alloys. For instance, advanced high-strength steels (AHSS)and “GigaPascal steels” have tensile strength in excess of 780 MPa and 1GPa, respectively. Matlock and Speer predicted that the strength of the3rd generation AHSS could potentially approach ˜2 GPa. In contrast, itremains a tremendous challenge for Al alloys to achieve a flow stress of1 GPa.

The best commercial Al alloys (typically age hardened) have a yieldstrength of ˜0.7 GPa. High strength (˜1 GPa) has hitherto beenachievable in rare cases, in some transition metal and/or rare earthelement-incorporated Al based metallic glasses and ceramicparticles-embedded composites. However, Al based metallic glassestypically show poor ductility. Grain refinement via severe plasticdeformation or cryomilling can increase the tensile strength of Alalloys through confinement of the dislocation migration. However, graingrowth occurs at ambient temperature and under low flow stress fornanocrystalline (nc) metals. Moreover, significant reduction of grainsize in nc monolithic and alloyed metals have shown softening phenomenon(smaller grains lead to lower yield strength), resulting from grainboundary-mediated activities when grain size falls below critical value.

Thus, there exists an unmet need for fabricating high-strength aluminumalloys and/or high-strength al-alloy coatings without compromisingductility and without any attendant softening phenomenon observed innanocrystalline metals and alloys.

SUMMARY

A high-strength aluminum alloy coating is disclosed. The coatingincludes aluminum 9R phase, fine grains, nanotwins, stacking faults, anda solute capable of stabilizing the 9R phase, the fine grains, and thestacking faults.

A method of making a high-strength aluminum alloy coating on a substratedisclosed. The method includes, providing a substrate, depositing atomsof the constituents of an aluminum alloy on the substrate utilizing adeposition method, such that the deposited atoms form a high-strengthaluminum alloy coating containing 9R phase, fine grains, nanotwins, andstacking faults.

A high-strength deformation layer in and on a casting of an aluminumalloy is disclosed. The high-strength deformation layer includes analuminum alloy comprising aluminum, 9R phase, fine grains, nanotwins,stacking faults, and a solute capable of stabilizing the PR phase, thefine grains, and the stacking faults.

A method of making a high-strength deformation layer in and on a castingof an aluminum alloy is disclosed. The method includes providing acasting of an aluminum alloy, and deforming the alloy by a deformationmethod, such that the deformation results in a high-strength aluminumalloy comprising a deformation layer containing 9R phase, fine grains,nanotwins, and stacking faults.

BRIEF DESCRIPTION OF DRAWINGS

Some of the figures shown herein may include dimensions. Further, someof the figures shown herein may have been created from scaled drawingsor from photographs that are scalable. It is understood that suchdimensions or the relative scaling within a figure are by way ofexample, and not to be construed as limiting.

FIG. 1A is an image of bright-field plan-view of monolithic Al (111)film.

FIG. 1B is an XTEM image of a cross-section of monolithic Al (111) film.

FIG. 1C shows (111) x-ray pole figure of monolithic Al indicating highlytextured (111) Al.

FIG. 1D is a plan-view TEM image of sputtered Al-2.5% Fe films withnanocolumns. The inserted SAED pattern displays slight in-plane rotationamongst columnar grains.

FIG. 1E is a cross-section TEM image of sputtered Al-2.5% Fe films withnanocolumns. The inserted SAED pattern shows the formation of twins.

FIG. 1F is the (111) pole figure of Al-2.5% Fe specimen. The six brightspots with similar intensity in the (111) pole figure demonstrates ahighly twinned Al-2.5% Fe specimen with 6 fold symmetry. The averagecolumnar grain size of Al—Fe films declines sharply with increasing Feconcentration.

FIG. 1G shows plots of hardness and average grain column size of Al—Fealloys of this disclosure as a function of iron composition.

FIG. 1H shows the Hall-Petch plot for the Al-xFe specimens and theselected ufg and nc monolithic Al and Al-xFe alloys processed viadifferent techniques. In comparison to the monolithic Al and Al-xFealloys, the sputtered Al—Fe films have exceptionally high flow stress(hardness/2.7).

FIG. 1I shows Ashby map of specific strength vs. specific modulus forvarious alloys.

FIG. 2A shows a low magnification cross-section TEM image ofas-deposited Al-2.5% Fe film showing columnar grains, ˜42 nm indiameter; the inserted SAED pattern shows typical twin structures inhighly (111) textured films.

FIG. 2B shows ae magnified TEM micrograph of 3 adjacent columns inas-deposited Al-2.5% Fe films and the corresponding fast Fouriertransform (FFT) patterns indicate column I and III have twin relation,whereas the FFT of column II shows the formation of 9R structures.

FIG. 2C illustrates the formation the microstructure of Al-2.5% Fefilms.

FIGS. 2D through 2F show HRTEM micrographs of the three boxes in FIG. 2Bshowing that column II is fully populated with the 9R phase.

FIGS. 2G through 2H show magnified HRTEM micrographs confirming theformation of intermingled 9R phase across the phase boundary (PB).

FIGS. 3A through 3C are SEM micrographs of pillars of Al(111) (FIG. 3A)and Al-xFe (x=2.5 or 5.9 at. %, FIGS. 3B and 3C respectively) capturedduring in situ pillar compression tests at different strains.

FIG. 3D shows true stress-true strain curves from compression tests onpillars fabricated from as-deposited Al(111) and Al—Fe films.

FIG. 3E shows the flow stresses and hardening exponent vs. Fe contentderived from compression tests on pillars fabricated from as-depositedAl(111) and Al—Fe films.

FIGS. 4A through 4C show MD simulations of uniaxial compressions on Al-5at. % Fe with certain columns fully populated with 9R phase bounded bytwo perfect crystal columns with twin relation (T and M denote twin andmatrix, respectively).

FIG. 4D shows a plot of compressive stress v. strain for alloy Al+5atomic % Fe with 9R phase.

FIG. 4E shows post-mortem bright-field and corresponding (200)_(T) and(200)_(M) dark-field TEM images of the deformed Al-2.5% Fe pillar.

FIG. 4F shows TEM image demonstrating 9R phase that survives thedeformation is greatly disturbed after the deformation.

FIG. 5A shows a digital micrograph of ˜20 μm thick Al-6 at. % Fe coatingthat well adheres to Si wafer with 3 inches in diameter.

FIG. 5B shows a cross-sectional SEM micrograph of the 20 μm thick Al-6at. % Fe coating.

FIG. 5C shows corresponding EDS patterns of Al and Fe and Si.

FIG. 5D shows the In situ compression tests applied on micropillars ofAl-6 at. % Fe coating.

FIG. 5E shows the stress vs. strain curves for micropillars of Al-6 at.% Fe coating.

DETAILED DESCRIPTION

For the purposes of promoting an understanding of the principles of thedisclosure, reference will now be made to the embodiments illustrated inthe drawings and specific language will be used to describe the same. Itwill nevertheless be understood that no limitation of the scope of thedisclosure is thereby intended.

In this disclosure, the words film and coating are used interchangeablyand mean a surface layer formed on a physical object. Also, it is to benoted that in structures containing nanotwins (known to those skilled inthe art), there is inter-twin spacing associated with the nanotwins. Theinter twin spacing is similar to grain size or column size that variesfrom ˜120 nm to ˜2 nm with different Fe compositions.

Grain Boundary (GB) engineering is an important avenue to improving themechanical strength of materials as investigated experimentally and bymolecular dynamic (MD) simulations in literature. An effective approachto accomplish high strength via GB engineering is to use low-Σboundaries, particularly Σ3 {111} coherent twin boundaries (CTBs). Therelation between the number of lattice points in the unit cell of acoincident site lattice (CSL) and the number of lattice points in a unitcell of the generating lattice is called Σ (Sigma); Sigma is the unitcell volume of the CSL in units of the unit cell volume of theelementary cells of the crystals. These notations and aspects are welldescribed in literature and are well understood by those skilled in inthe area of materials engineering. Nanotwinned (nt) metals with CTBsexhibit high strength, ductility, electrical conductivity, and superiorthermal stability. CTBs serve as an effective strengthening motif byblocking dislocation motion and can also accommodate plastic strain bystoring dislocations and even migrate under stress. Prior works on ntmetals mostly concentrate on metals with low-to-intermediatestacking-fault energy (SFE, γ_(sf)), such as Ag (SFE=16 mJ/m²), Cu(SFE=45 mJ/m²), Cu alloys, and 330 austenitic stainless steels(SFE=10-20 mJ/m²), due to the ease of formation of growth twins. In Alwith high SFE of 120-144 mJ/m², sporadic deformation twins have beenobserved in deformed nanocrystals or under extreme conditions, e.g. nearcrack tips or indentation edge, under high stress and strain rate ordeformation at cryogenic temperature. Although dislocation-twininteractions in twinned Al have been studied by MD simulations, itremains challenging to introduce high-density growth twins in Al and theconcept of engineering twin boundaries (TBs) to macroscopicallystrengthen Al and Al alloys has not been actualized. Recently,incoherent twin boundaries (ITBs) have been successfully introduced intoAl by adopting nt Ag as seed layers. In situ nanoindentation and MDsimulations prove that the ITBs in Al act as strong barriers todislocation pile-up, and eventually allow dislocation transmission atthe steps along ITBs. The hardness of these Al films remains low, ˜1GPa, corresponding to a flow stress of 300 MPa.

In this disclosure, the fabrication of high-strength, twinned Al-xFe(x=1-10 at. %) solid solution films with a substantial volume fraction(˜25%) of 9R phase is described. 9R phase is well understood by thoseskilled in the art. In this disclosure, in describing Al—Fe alloys, % Ferefers to atomic percent of iron. The hardness of Al—Fe films of thisdisclosure reaches 5.5 GPa. In situ pillar compression tests reveal thatthe Al—Fe films of this disclosure have flow stress exceeding 1.5 GPa,comparable to high-strength martensitic steels, and exhibit giant strainhardening ability. DFT (density function theory) calculations show thatthe addition of Fe increases the SFE of Al. However, Fe increases theratio between unstable SFE and stable SFE (γ_(us)f/γ_(sf)) andsignificantly enhances the stability of growth twins and 9R phase. MDsimulations show that the ultra-high strength of Al—Fe films of thisdisclosure derives from the dislocation-ITB and 9R phase interactions.

In experiments leading to this disclosure, Al and Al-xFe (x=1-10 at. %)alloys were deposited onto single crystal Si(111) substrates by directcurrent magnetron sputtering at a base pressure of 6-10×10⁻⁸ Torr usingAl (99.99%) and Fe (99.95%) targets. The co-sputtering of Al—Fe alloyscan scale up the coatings in thickness by increasing deposition time andtwo exemplary thicknesses of coatings herein demonstrated are 1 μm and20 μm. Plan-view and cross-section transmission electron microscopy(TEM) specimens were prepared by mechanical grinding and dimpling,followed by low-energy ion milling/polishing, and examined in FEI TecnaiG2 F20 and Talos 200× microscopes operated at 200 kV, equipped withFischione ultra-high resolution high angle annular dark field (HADDF)detectors. X-ray diffraction (XRD) patterns were acquired on aPanalytical Empyrean X'pert PRO MRD diffractometer with a Cu Kai source.Pole figure measurements were conducted with the tube-in-line focus anda capillary X-ray lens with a spot opening of 1×1 mm² on the primaryoptics and parallel plate collimator on the detector side. FIGS. 1A and1B are plan-view and cross-section TEM images indicating submicroncolumnar grains of monolithic Al grown on Si (111) substrate. Theinserts in FIGS. 1A and 1B show selected area electron diffraction(SAED) patterns indicating that {111} Al is grown epitaxially on Si(111)substrates. FIG. 1C shows (111) x-ray pole figure of monolithic Alindicating highly textured (111) Al. FIG. 1D is a plan-view TEM image ofsputtered Al-2.5% Fe films with nanoscale columnar grains (referred toas nanocolumns). The inserted SAED pattern displays slight in-planerotation amongst columnar grains. FIG. 1E is a cross-section TEM imageof sputtered Al-2.5% Fe films with nanocolumns. The SAED pattern in theinsert of FIG. 1E shows the formation of twins. FIG. 1F is the (111)pole figure of Al-2.5% Fe specimen. The six bright spots with similarintensity in the (111) pole figure demonstrates a highly twinned Al-2.5%Fe specimen with 6 fold symmetry. The average columnar grain size ofAl—Fe films declines sharply with increasing Fe concentration. Al grownepitaxially on Si(111) substrates with an average grain size of ˜450 nm.In contrast, columnar grain size decreases to ˜40 nm in Al-2.5% Fe(FIGS. 1D through 1E) and ˜4 nm in Al-5.9% Fe as seen in theexperiments. The SAED pattern of the plan-view TEM micrograph of Al-2.5%Fe also displays (111) texture, consistent with XRD patterns obtained.Referring to FIG. 1E, The SAED pattern of the XTEM (meaningcross-sectional TEM) micrograph of the Al-2.5% Fe specimen (FIG. 1E)reveals the classical twinned diffraction pattern examined along the Al<011> zone axis. Comparisons of XRD (111) pole figures show thatmonolithic Al has 3 fold symmetry, whereas a six-fold symmetry ofequally bright {111} poles emerges in Al-2.5% Fe, indicative of a largefraction of twin variants in the highly textured films. Similarly, thetwinned patterns were observed in XRD pole figure analyses along (200)and (220) poles in the studies leading to this disclosure.

Hardness and modulus of the alloys were determined using theinstrumented nanoindentation technique on both Fischerscope 2000XYpnano/micro-indenter and Hysitron TI950 at different indentation depths.Maximum indentation depth was controlled not to exceed 15% of the filmthickness to avoid substrate effect. FIG. 1G shows plots of hardness andaverage grain column size of Al—Fe alloy coatings with a thickness of 1μm as a function of iron composition. The hardness of sputtered Al-5.9%Fe specimens approaches 5.6 GPa, much greater than that of monolithic ntAl, ˜0.7 GPa, and previously reported ufg, nc Al and Al—Fe alloys. Thecolumnar grain size decreases from ˜130 to 2 nm with increasing Feconcentration. The evolution of hardness of nt Al—Fe measured bynanoindentation technique with composition of Fe is shown in FIG. 1G.The hardness of the nt Al alloys starts at 2.5±0.15 GPa for Al-1.1% Fe,reaches a maximum value of 5.5±0.1 GPa for Al-5.9% Fe, and thendecreases to 4.7±0.2 GPa for Al-9.9% Fe. It has been reported inliterature that a quaternary Al₈₄Ni₇Gd₆Co₃ alloy containing nanoscaleAl₁₉Gd₃Ni₅ and Al₉Co₂—Al₃Gd intermetallics (with merely 23% fcc Al)fabricated by hot pressing of powders after ball milling for 100 h hasshown similar level of hardness. First principle density function theory(DFT) calculations predicted virtually zero strength for Al—Fe withparts per million of Fe solutes. Prior studies on Al—Fe alloys aresummarized in FIG. 1G. The current nt Al—Fe alloys have the higheststrength among all Al—Fe alloys reported to date. Epitaxial nt Al with˜200 nm ITB spacing has a hardness of 1.2 GPa, comparable to thehardness of nc Al with a grain size of 30-50 nm. Referring to FIG. 1G,the hardness of sputtered Al-5.9% Fe specimens approaches 5.6 GPa, muchgreater than that of monolithic nt Al, ˜0.7 GPa, and previously reportedultra-fine grained (ufg), nc Al and Al—Fe alloys. The columnar grainsize decreases from ˜130 to 2 nm with increasing Fe concentration.

The mechanical performance of nt Al—Fe is further analyzed by using flowstress converted from hardness divided by a Tabor factor of 2.7 in aHall-Petch plot to investigate size-dependent strengthening. FIG. 1Hshows the Hall-Petch plot for the Al-xFe specimens and the selected ufgand nc monolithic Al and Al-xFe alloys processed susvia differenttechniques. In comparison to the monolithic Al and Al-xFe alloys, thesputtered Al—Fe films have exceptionally high flow stress(hardness/2.7). A variety of solutes, such as Ag, Mo, Cr and W, wereselected to alloy with Al using similar deposition parameters, but thehardness of these binary Al alloys hardly exceeded 3.5 GPa while keepingsolute concentration to 10 at. % or less, as evidenced from the studiesleading to this disclosure.

FIG. 1I show the Ashby map, which refers to a map showing specificstrength vs. specific modulus for various alloys. In general, materialswith high specific strength and high specific modulus are preferred. Inother words, it is desirable for the materials whose values fall intothe upper right portion of the map. Conventional Al alloys already havehigh specific modulus, but they have low specific strength as shown inFIG. 1 . An objective of this disclosure is to significantly boost thespecific strength of the Al alloys. The high-strength aluminum coatingsof this disclosure appear in the upper right portion of the plot in FIG.1I. To facilitate the TEM sample preparations, the film thickness inFIG. 1 is 1 μm.

FIGS. 2A through 2H show microstructures of as-deposited Al-2.5% Fefilms. FIG. 2A shows a low magnification cross-section TEM image ofas-deposited Al-2.5% Fe film showing columnar grains, ˜41 nm indiameter; the insert of FIG. 2A, the SAED pattern, shows typical twinstructures in highly (111) textured films. FIG. 2B shows magnified TEMmicrograph of 3 adjacent columns, indicated as I, II, and III in FIG.2B, in as-deposited Al-2.5% Fe film; the corresponding fast Fouriertransform (FFT) patterns indicate column I and III have twin relation,whereas the FFT of column II shows the formation of 9R structures. FIG.2C is a schematic illustration of the formation the microstructure ofAl-2.5% Fe films. FIGS. 2D through 2F show High-resolution TEM (HRTEM)micrographs of the three boxes, labeled d, e, and f, in FIG. 2B showingthat column II is fully populated with the 9R phase. FIGS. 2G through 2Hshow magnified HRTEM micrographs confirming the formation ofintermingled 9R phase across the phase boundary (PB). Intermingling hererefers to intermixing of two 9R phases. Referring to FIG. 2G, themagnified HRTEM micrographs highlight the diffused left PB of the 9Rphase containing intermingled structure (repeating SFs in 9R at twosides of PB have one {111} interplanar distance difference) and therelatively sharp right PB.

An examination of XTEM micrograph of the Al-2.5% Fe films in FIG. 2Ashows curvy grain boundary geometry and, interestingly 9R phase entirecolumns in many cases (FIG. 2B)), and the giant 9R (˜50 nm in width) isoften bounded by two adjacent columns with twin relations as shown bythe fast Fourier transform (FFT) as insets (FIG. 2B). The volumefraction of 9R phase is estimated as ˜25%, and the microstructure of thefilm containing 9R phase is shown schematically in FIG. 2C. Highresolution TEM (HRTEM) micrographs, FIGS. 2D through 2F show the left,middle and right sections, respectively, of a typical 9R structure. The9R phase typically contains three Shockley partials, one edge and twomixed partials, on adjacent {111} planes. An intermingled 9R structureis found on the left boundary of the 9R phase (FIG. 2D). Few regionsthat ensemble perfect crystal are also observed and break the continuityof 9R phase (FIG. 2E).

To investigate the deformability of the Al and Al—Fe alloys, pillarswith ˜500 nm in diameter and a diameter-to-height aspect ratio of ˜1:2were fabricated through focused ion beam (FIB) and a series ofconcentric annular milling and polishing with progressively de-escalatedcurrents were applied to minimize the substrate engagement and tapering.A Hysitron PI 87xR PicoIndenter equipped with a 5 μm diamond flat punchtip was used to conduct in situ compression experiments inside an FEIquanta 3D FEG scanning electron microscope. A piezoelectric actuator onthe capacitive transducer enables the collection of force-displacementdata and the morphological evolution of the micropillar was capturedconcurrently during real-time deformation. An average drift rate of 0.2nm/s was estimated in preloading process and all the experiments werefinished in about 60 s and the estimated noise for force measurement is±5 μN. The ex-situ compressions were performed using 8 μm flat punch inHysitron TI 950 TriboIndenter equipped with a transducer with 1 nN loadresolution and the thermal drift rate was below 0.05 nm/s and monitoredfor 40 s prior to compressions conducted at a strain rate of 1×10⁻³/sfor ˜150 s.

In situ uniaxial compression tests were performed on pillars fabricatedfrom as-deposited Al(111) and Al—Fe films performed inside a scanningelectron microscope. FIGS. 3A through 3C show SEM micrographs of pillarsof Al(111) and Al-xFe (x=2.5 or 5.9 at. %) captured during in situpillar compression tests at different strains. The Al—Fe specimensexhibit ductile deformation behaviors with no detectable shear banding(offset) event. Referring to FIGS. 3 B and 3 C, it can be seen that thedeformed Al—Fe pillars experience dilation on the top, in contrast tothe conventional plastic barreling observed in monolithic nt Al as wellas conventional softening. True stress-strain curves of the pillarsshown in FIG. 3D indicate that the flow stress of Al-2.5% Fe and Al-5.9%Fe exceeds 1 and 1.6 GPa, respectively. Partial unloading experiment wasperformed to measure the elastic modulus of the pillars. The truestress-strain curves of the deformation volume, i.e. the dilated volume,is estimated by analyzing the instantaneous geometry variationsynchronized to the mechanical response, the detailed methodology ofwhich is explained in SI. The flow stresses at ˜10% strain, shown inFIG. 3E, shows that the Al—Fe alloy films have significantly higher flowstress than monolithic Al films. Also, from FIG. 3E, the strainhardening exponent, n, of the Al—Fe films are found to be greater thanthat of monolithic Al films.

Further observations on the deformation experiments: In situ SEMmicrographs in FIG. 3A shows ductile deformation of monolithic Al, asevidenced by barreling of the pillar, and the specimen is very soft,with a flow stress of 200 MPa (FIG. 3D). The SEM snapshots of Al-2.5% Fealloys taken during deformation up to strain of 20% (FIG. 3B) showextensive deformation of the pillar, with a reverse conicalfrustum-shaped geometry in contrast to barreling of the compressed Alpillars. The calculated flow stress of Al-2.5% Fe after considering thegeometry of pillars exceeds 1 GPa. Similar geometry change was observedin the Al-5.9% Fe micropillars (FIG. 3C). No shear offsets were observedin the specimen. At 10% strain, the flow stress of the Al-5.9% Feexceeds 1.5 GPa. The ultra-high strength Al-5.9% Fe pillars can toleratemore than 50% true strain with no cracks. In addition to monotonicloading, multiple tests with numerous partial unloading were performed(FIG. 3D) to check the modulus and the reliability of methodology. Themeasured modulus of pillars ranges from 85 to 115 GPa.

Al is known to be soft and thus various types of strengtheningapproaches have been applied. Among these methods, age hardening throughprecipitation is widely used. Solid solution hardening alone isgenerally considered an inefficient method to accomplish significantstrengthening in Al alloys. Grain refinement has also been widelyadopted to increase the strength of Al alloys. But to date, the bestcommercial Al alloys, including 7XXX series Al alloys, only have flowstress up to 700 MPa. Tempering of certain 7XXX series of Al alloys hasled to a maximum flow stress of ˜1 GPa. For the coatings of thisdisclosure, the hardness and flow stress of nt Al—Fe reach 5.5 GPa and1.5 GPa, respectively. The nt Al—Fe alloys exhibit superior strengthcompared with many other alloys on the Ashby map of specific strengthversus specific elastic modulus (FIG. 1I). Three strengtheningmechanisms, i.e. grain boundary, solid solution and dislocationstrengthening, might contribute to the anomalously high strength. First,it has been widely adopted that nano-grains explain the improvedstrength through the constraint of dislocation motions, as elucidated bythe empirical Hall-Petch law, but the strength of nanocrystalline Alhardly exceeds 600 MPa and softening may occur (smaller grain size leadto lower hardness) below certain grain sizes based ondislocation-accommodated grain boundary sliding. Second, the solidsolution hardening estimated by the Fleischer equation is merely 0.26GPa for Al-6% Fe.

High density of stacking faults in form of 9R phase as preexistingdefects might contribute to significant strengthening and thesolute-decorated ITBs may retard ITB migration and restrict dislocationtransmission. The role of 9R on strengthening of metallic materials islargely overlooked. This is because that 9R phase is rare, even inmetals with low SFE, such as Cu and Ag. In general, high densitynanotwins are considered significant in achieving high strength and goodductility in nt metals with low SFEs. 9R phase is scarce in Al due toits high SFE. Even under high shear stress 9R phase has not beenobserved in Al or its alloys to date, in sharp contrast to low SFE Cuwhere 9R could form and propagate under shear stress. An asymmetric tilton the {111} planes of 9R phase is detected relative to its two boundingcolumns. The breakage of the symmetry could be ascribed to the energyminimization by structural relaxation envisioned by MD and the 9R widthis a function of inclination angle for asymmetric tilt GB s. Themisalignment amongst nanoscale columns in Al—Fe may facilitate or resultfrom the 9R formation.

Recently ITBs have been successfully introduced into monolithic Al usingAg template. The Ag seed layer contains high-density growth twins due toits ultra-low SFE, 8 mJ/m². The TBs, mostly ITBs formed in Ag canpropagate readily into Al as Al and Ag have nearly identical latticeparameter (<1% lattice mismatch). In situ nanoindentation studies showthat ITBs in Al can induce high flow stress, ˜300 MPa.

An approach to induce high density TBs in Al would be to reduce its SFEby alloying as lower SFE can promote the formation of growth twins.Hence the observation of high-density 9R phase in Al—Fe alloy suggeststhat the SFE of Al might have been reduced. However, DFT calculations instudies leading to this disclosure show just the opposite trend, thatis, the SFE of Al—Fe is in fact higher than that of Al. Hence it isnatural to ask why 9R phase would still form in Al—Fe alloy if its SFEis even greater than Al? The answer, without being bound by theory, isconsidered to be the significantly enhanced stability of 9R phase inAl—Fe due to the addition of Fe solutes. The equilibrium SF spacing inpure Al is small, ˜0.3 nm, and the formation of SFs will naturallyundergo a de-faulting process. De-faulting, in the context of thisdisclosure is intended to mean reduction or elimination of stackingfaults or undoing the stacking faults. Similarly, the formation ofgrowth and deformation twins in Al is also difficult, though feasible(when grain size is in the nanoscale regime). MD simulations in thestudies leading to this disclosure show that TBs or SFs artificiallyintroduced into Al will automatically de-twin (undoing the twins) orde-fault at room temperature. In contrast, high density SFs and TBsinduced in Al—Fe are rather stable at room temperature. It is worthmentioning that Fe atoms are uniformly distributed throughout the entirefilms as examined by atomic resolution STEM.

MD simulations were performed to compare the mechanical behavior of ntAl with ITBs and nt Al—Fe with high-density SFs and 9R phase. Bothmaterials have identical columnar grain sizes, 20 nm. In general, ITBsare strong barriers to the transmission of dislocations, and the yieldstrength and peak strength of nt Al reaches ˜2 and 4 GPa, respectively.In nt Al—Fe with high density 9R phase, although the yield strength iscomparable to that of nt Al, a much higher work hardening rate isobserved, leading a peak strength of nearly 6 GPa. Significant workhardening observed in nt Al—Fe during MD simulations is in goodagreement with the in situ compression results. MD simulationsdemonstrated that high-density stacking faults in Al—Fe are responsiblefor strong flow stress oscillation compared to the pure Al samplewithout stacking faults. Stacking faults impede the glide ofdislocations and thus strengthen the Al alloy (corresponding to strainhardening). FIGS. 4A through 4C show MD simulations of uniaxialcompressions on Al-5 at. % Fe with certain columns fully populated with9R phase bounded by two perfect crystal columns with twin relation (Tand M denote twin and matrix, respectively). The Fe solutes areuniformly distributed through the alloy films and the 9R columns andperfect crystal columns are indicated by arrows. The as-constructedstructure and deformed structure at 8 and 10% strains, respectively, aredirectly compared in both 3D and 2D modes. MD studies show Fe-inducedstability and sectional recovery of 9R phase as well as dislocation-9Rinteractions in Al-5% Fe specimen. Stacking faults in monolithic Alundergo recovery to perfect crystal at ambient temperature, whereas aconsiderable fraction of 9R phase survives deformation. Meanwhilepartial 9R phase recovers, but left debris. The increased SFE makesdefaulting/detwinning even more difficult. The dislocations are blockedby the stable and thick 9R and the strain hardening throughdislocation-9R interactions is further promoted. Furthermore, the 9R isdeformable and high-density partials within the 9R phase can migrate toaccommodate plasticity. FIG. 4D shows a plot of compressive stress v.strain for alloy Al+5 atomic % Fe with 9R phase. FIG. 4E showspost-mortem bright-field and corresponding (200)_(T) and (200)_(M)dark-field TEM images of the deformed Al-2.5% Fe pillar. FIG. 4F showsTEM image demonstrating 9R phase that survives the deformation isgreatly disturbed after the deformation. Referring to FIG. 4F, Distorted9R is bounded by tilted 9R and undisturbed 9R.

The work hardening exponent in nt Al—Fe (shown in FIG. 3E) measured byin situ compression tests is ˜0.6, comparing to ˜0.4 for monolithic Alsubjected to micropillar compression tests at the same strain rate. Inaddition, substantial load drop was observed in monolithic Al in FIG.4D, whereas nt Al—Fe does not show any load drop during compressiontests. In 3D MD simulations (FIGS. 4A through 4C), the substantial workhardening in nt Al—Fe can be attributed to extensive interactions ofdislocations with SFs and 9R phase. Similarly, 2D MD simulations on amuch greater cell, containing 9R phase in alternating grains (similar toexperimental observations) show that the interaction of glidedislocations with 9R phase leads to work hardening. The high ductilityof nt Al—Fe is thus underpinned by 9R phase dominated work hardeningduring compressive deformation. Of course, the ultra-high strengthaccomplished in Al—Fe is also related to the nanoscale grain size. Butgrain size alone is insufficient to explain the high strength in thematerials.

It should be noted that all solutes do not have the same capability tostabilize the 9R phase in Al, and their effectiveness varies. Forexample, solutes such as molybdenum and tungsten are less effective toreduce the average grain size of Al comparing to Fe solute and hencethese solutes may only moderately increase the strength of Al alloys.This inference is based on extensive studies performed throughexperiments leading to this disclosure on numerous other binary Alalloys, by using, Mo, W, Ag, and Cr as solutes. The studies showed thatthat these solutes were not as effective as Fe, Ni and Co to reducegrain size and/or produce 9R, and hence led to relatively insignificantstrengthening effect. However, it may be possible to tailor processingconditions and chemistry to tune the effectiveness of different solutesin influencing grain size. Further, it should be recognized that therecan be other solutes that can promote the formation of 9R phase andnanograins may also lead to high strength and plasticity in Al alloys.

It has been stated that FCC metals will undergo 9R expansion wheredouble core Shockley partials pin one side of ITBs while shear stress isapplied along <111> direction but experience boundary sliding and ITBand Shockley partial decoupling at applied shear along <110> direction.The texture of nt Al—Fe specimens may benefit the superb plasticity.

In this disclosure, high density twins and 9R phase are introduced intoAl by using Fe solutes. The nt Al—Fe alloys achieve a hardness and aflow stress as high as 5.5 GPa and ˜1.5 GPa, respectively, which iscomparable to high strength martensitic steels. The in situ compressionsreveal a ductile deformation of the nt Al—Fe alloys as well ascapability for tremendous strain hardening. DFT calculations show thatthe Fe addition counterintuitively increases the SFE of Al so that ahigher energy penalty is required to undergo detwinning and defaultingprocesses for the twin and 9R phase. MD simulations suggest the additionof Fe can appreciably improve the stability of 9R phase for furtherstrain hardening the materials and the dislocation-ITB s interactionalso attribute to the high strength and giant strain hardening.

It should be noted that based on fabrication of high-strength Al alloycoatings, bulk nanotwinned high-strength Al alloys can be fabricated. Anon-limiting method to achieve this includes applying a surfacemechanical grinding treatment technique to an Al ally casting tofabricate nt Al alloys. This severe plastic deformation technique canallow fabrication of 100 micron thick nanocrystalline Al alloy plates.The average grain size can approach 50-200 nm. Deformation willintroduce high-density twins and stacking faults, and 9R phase. As thesolutes will stabilize these stacking faults, the as-processed Al alloyscan have high strength, 0.7-2 GPa, and high hardness, 2-5 GPa in thehighly deformed surface layer of Al alloys.

Another method of achieving a high-strength Al alloy sheets includesapplying rolling and forging to significantly reduce the thickness of Alalloy plates. Such dramatic rolling and forging experiments will alsoinduce nanograins, high-density nanotwins and stacking faults into ourAl alloys. This technique can provide Al alloy sheets that have 1-5 mmin thickness. By extensive application of such deformation processes,bulk nanotwinned Al alloys can be fabricated.

To evaluate the mechanical properties of thick Al—Fe coatings, 20μm-thick Al—Fe coatings were deposited on Si wafers with ˜3 inch indiameter using the magnetron sputtering technique with the samedeposition parameters. FIG. 5A shows a digital micrograph of ˜20 μmthick Al-6 at. % Fe coating that well adheres to Si wafer with 3 inchesin diameter. FIG. 5B shows a cross-sectional SEM micrograph of the 20 μmthick Al-6 at. % Fe coating and FIG. 5C shows corresponding EDS patternsof Al and Fe and Si. FIG. 5D shows the In situ compression tests appliedon micropillars of Al-6 at. % Fe coating. FIG. 5E shows the stress vs.strain curves for these micropillars with 3 μm in diameter. Thestress-strain curve of FIG. 5E is in good agreement with resultsobtained from micropillars with 500 nm in diameter and ˜1 μm in heightpresented above. These micropillars of large diameters showexceptionally high flow stress (>1.5 GPa) and strain hardening ability.Prior studies show that extrinsic size effect (“smaller is stronger”)takes place when pillar diameter falls below critical value comparableto or finer than intrinsic microstructural features, where strengtheningmay arise from dislocation starvation for nanopillars made of single orcoarse-grained crystals. To avoid such an extrinsic size effect inducedstrengthening, pillars with 3 μm in diameter and 7.5 μm in height havebeen fabricated and tested under the same strain rate. The flow stressof these large pillars remains high, exceeding 1.5 GPa (FIG. 5E),similar to the mechanical behaviors of pillars with 500 nm in diameter.

Based on the above detailed description, it is an objective of thisdisclosure to describe a high-strength aluminum alloy coating on acasting or a substrate. In one non-limiting embodiment, the castingcomprises Al-5% Fe. The high-strength aluminum alloy coating of thisdisclosure contains aluminum, 9R phase, fine grains, nanotwins, stackingfaults, and a solute capable of stabilizing the 9R phase, the finegrains, and the stacking faults. The nanosized grains with Fe-stabilized9R, predominately account for the strength improvement. Fe enables thenanograin formation and the stability of 9R. To ensure that there ishigh strength, it is desirable at least nanocrystalline grain size. But9R phase will make nanocrystalline metals even stronger. And moreimportantly, 9R phase enable outstanding deformability (plasticity).Thus in some embodiments of the coatings of this disclosure one or moreof the listed features such as for example purposes only, 9R phase,nanotwins and stacking faults, may be absent, and yet some strengthimprovement can be realized.

The thickness of the high-strength aluminum alloy of this coating can bein the range of 0.1-200 micrometers. In this high-strength aluminumalloy coating the fine grains can be equiaxed or columnar. Anon-limiting range for the size range of the fine grains in thesecoatings is 2 nm-10,000 nm. In some preferred embodiments of thecoatings of this disclosure, the fine grains are in the size range of 5nm-200 nm. In the high-strength aluminum alloy coatings of thisdisclosure, the inter-twin spacing is in the range 1 nm-200 nm. Examplesof solutes in the high-strength aluminum alloy coatings of thisdisclosure include, but not limited to, iron, cobalt, titanium,magnesium, and chromium. In some preferred embodiments of the coatingsof this disclosure, the solute is iron and the iron content is in therange of 1-30 atomic percent.

In some embodiments of high-strength aluminum coating of thisdisclosure, the flow stress of the coating is in the range of 0.7-3.0GPa and ductility of the coating is in the range of 5-30%. In someembodiments of the high-strength aluminum coating of claim 16, the flowstress of the coating is in the range of 1.0-2.0 GPa and ductility ofthe coating is in the range of 10-25%. In some embodiments of thehigh-strength aluminum coating of this disclosure, the hardness of thecoating is in the range of 2.0-9.0 GPa and ductility of the coating isin the range of 5-30%. In some embodiments of the high-strength aluminumcoating of this disclosure, the hardness of the coating is in the rangeof 2.0-9.0 GPa and ductility of the coating is in the range of 10-25%.

It is another objective of this disclosure to describe a method ofmaking a high-strength aluminum alloy coating on a substrate. The methodincludes depositing on a substrate, atoms of the constituents of analuminum alloy on the substrate utilizing a deposition method, whereinthe deposited atoms form a high-strength aluminum alloy coatingcontaining 9R phase, fine grains, nanotwins, and stacking faults.Non-limiting examples of alloying elements for the aluminum alloycoating formed in this method include iron, cobalt, titanium, magnesium,and chromium. The deposition methods include, but not limited tosputtering, evaporation, laser ablation, and physical vapor deposition.The substrate on which the high-strength aluminum-alloy coating isformed can be a semiconductor. Non-limiting examples of semiconductorssuitable for this method include silicon, germanium, and galliumarsenide. In some embodiments of this method, the substrate can be ametal or an alloy. Non-limiting examples of metals suitable assubstrates for this method include copper, nickel, aluminum, iron, andtitanium. Non-limiting examples of alloys suitable as substrates forthis method include an aluminum alloy copper alloy, a nickel alloy and atitanium alloy. In some embodiments of the method, the flow stress ofthe coating formed is in the range of 0.7-3.0 GPa and ductility of thecoating is in the range of 5-30%. In some other embodiments of themethod, the flow stress of the coating is in the range of 1.0-2.0 GPaand ductility of the coating is in the range of 10-25%. In someembodiments of the method the hardness of the coating is in the range of2.0-9.0 GPa and ductility of the coating is in the range of 5-30%. Insome other embodiments of the method, the hardness of the coating is inthe range of 2.0-9.0 GPa and ductility of the coating is in the range of10-25%. In some embodiments of the methods of this disclosure one ormore of the listed features such as for example purposes only, 9R phase,nanotwins and stacking faults, may be absent, and yet some strengthimprovement can be realized.

It is another objective of this disclosure to describe a high-strengthdeformation layer in and on a casting of an aluminum alloy. Thehigh-strength deformation layer of this disclosure includes an aluminumalloy containing aluminum, 9R phase, fine grains, nanotwins, stackingfaults, and a solute capable of stabilizing the 9R phase, the finegrains, and the stacking faults. The high-strength deformation layer ofthis disclosure can contain fine grains which are equiaxed or columnar.The high-strength deformation layer of this disclosure can have finegrains in the size range of 2 nm-10,000 nm, a preferred range being 5nm-200 nm. In some embodiments of the high-strength deformation layer ofthis disclosure, inter-twin spacing is in the range 1 nm-200 nm.Examples of solute in the high-strength deformation layer of thisdisclosure include, but not limited to, iron, cobalt, titanium,magnesium, and chromium. In some preferred embodiments, the solute isiron in the range of 1-30 atomic percent. In some embodiments of thehigh-strength deformation layer of this disclosure one or more of thelisted features such as for example purposes only, 9R phase, nanotwinsand stacking faults, may be absent, and yet some strength improvementcan be realized.

It is yet another objective of this disclosure to describe a method ofmaking a high-strength deformation layer in and on a casting of analuminum alloy. The includes providing a casting of an aluminum alloy,and deforming the alloy by a deformation method, wherein the deformationresults in a high-strength aluminum alloy comprising a deformation layercontaining 9R phase, fine grains, nanotwins, and stacking faults. Thedeformation methods include, but not limited to grinding, rolling,extrusion, forging, and stamping. In some embodiments, the thicknessreduction of the casting due to deformation is in the range of 20-50%.Alloying elements for the aluminum alloy of the casting of the methodinclude, but not limited to iron, cobalt, titanium, magnesium, andchromium. In some embodiments of the method, the aluminum alloy of thecasting contains iron in the range of 1-30 atomic percent. In someembodiments of the method of making high-strength deformation layer ofthis disclosure one or more of the listed features such as for examplepurposes only, 9R phase, nanotwins and stacking faults, may be absent,and yet some strength improvement can be realized.

Another aspect of this disclosure is employing the methods offabricating the coatings of this disclosure to make articles andcomponents for use several applications. Other embodiments of thisdisclosure include articles comprising the Al—Fe coatings of thisdisclosure. These articles include, such as, but not limited to, Alalloy coatings for spatial light regulator, high-strength wear andscratch resistant coatings; and micro- and nanoelectromechanical systems(M/NEMS) device. These articles find applications in aerospace industry,protective coatings; optical devices, sensors and actuators.}

While the present disclosure has been described with reference tocertain embodiments, it will be apparent to those of ordinary skill inthe art that other embodiments and implementations are possible that arewithin the scope of the present disclosure without departing from thespirit and scope of the present disclosure. Thus, the implementationsshould not be limited to the particular limitations described. Otherimplementations may be possible. It is therefore intended that theforegoing detailed description be regarded as illustrative rather thanlimiting. Thus, this disclosure is limited only by the following claims.

The invention claimed is:
 1. A method of making a high-strength aluminumalloy coating on a substrate, the method comprising: providing asubstrate; depositing atoms of the constituents of an aluminum alloy onthe substrate utilizing a deposition method, wherein the deposited atomsform a high-strength aluminum alloy coating containing 9R phase, grainsin the size range of 2 nm-10,000 nm, nanotwins, and a solute capable ofstabilizing the 9R phase.
 2. The method of claim 1, wherein thedeposition method is one of sputtering, evaporation, laser ablation, andphysical vapor deposition.
 3. The method of claim 1, wherein thesubstrate is a semiconductor.
 4. The method of claim 3, wherein thesemiconductor is one of silicon, germanium, and gallium arsenide.
 5. Themethod of claim 1, wherein the substrate is a metal or an alloy.
 6. Themethod of claim 5, wherein the metal is one of copper, nickel, aluminum,iron, and titanium.
 7. The method of claim 5, wherein the alloy is oneof an aluminum alloy, a copper alloy a nickel alloy and a titaniumalloy.
 8. The method of claim 7, wherein the aluminum alloy comprisesone or more of iron, cobalt, titanium, magnesium, and chromium.
 9. Ahigh-strength deformation layer in and on a casting of an aluminumalloy, wherein the high-strength deformation layer includes an aluminumalloy comprising: aluminum; 9R phase; grains in the size range of 2nm-10,000 nm; nanotwins; and a solute capable of stabilizing the 9Rphase.
 10. The high-strength deformation layer of claim 9, wherein thefine grains are equiaxed or columnar.
 11. The high-strength deformationlayer of claim 9, wherein inter-twin spacing of the nanotwins is in therange 1 nm-200 nm.
 12. The high-strength deformation layer of claim 9,wherein the solute is one of iron, cobalt, titanium, magnesium, andchromium.
 13. The high-strength deformation layer of claim 12, whereinthe solute is iron.
 14. The high-strength deformation layer of claim 13,wherein the iron content is in the range of 1-30 atomic percent.
 15. Amethod of making a high-strength deformation layer in and on a castingof an aluminum alloy, the method comprising: providing a casting of analuminum alloy; and deforming the alloy by a deformation method, whereinthe deformation results in a high-strength aluminum alloy comprising adeformation layer 9R phase, grains in the size range of 2 nm-10,000 nm,nanotwins, and a solute capable of stabilizing the 9R phase.
 16. Themethod of claim 15, wherein the deformation method is one of rolling,extrusion, forging, and stamping.